Refractory composite alloys containing rod-like and/or platelet-like lamellae

ABSTRACT

Polyvariant alloys which comprise a matrix and a directional reinforcement phase imbedded in said matrix in form of chromiumfree mono-crystalline particles of carbide. The method for the manufacture of such alloys involves zoning a rod under determined conditions of movement rate and thermal gradient at the solidification interface of the rod.

United States Patent Bibring et al.

[ 1 Mar. 18, 1975 REFRACTORY COMPOSITE ALLOYS CONTAINING ROD-LIKE AND/ORPLATELET-LIKE LAMELLAE lnventors: Herv E. Bibring, Meudon; Georges P.Seihel, Cachan; Maurice Rabinovitch, Chatillon, all of France Assignee:Office National dEtudes et de Recherches Aerospatiales, Chatillon,France Filed: July 3, 1972 Appl. No.: 268,751

Related U.S. Application Data Continuation-impart of Ser. No. 2,160,Jan. 12,

1970, abandoned.

Foreign Application Priority Data Apr, 21, 1969 France 69.12452 Dec. 23,1969 France 69.44708 U.S. Cl. 29/l9l.2, 75/170 Int. Cl. 332g 15/00, C22C19/00 Field of Search 75/170 Primary Examiner-Allen B. Curtis Attorney,Agent, or Firm-Karl F. Ross; Herbert Dubno [57] ABSTRACT Polyvariantalloys which comprise a matrix and a directional reinforcement phaseimbedded in said matrix in form of chromium-free mono-crystallineparticles of carbide. The method for the manufacture of such a1- loysinvolves zoning a rod under determined conditions of movement rate andthermal gradient at the solidification interface of the rod.

16 Claims, 19 Drawing Figures PATENTEDMAR] 81975 SHEET (2F 10 ig-E I cs

HERVE "LBIBRING GEORGES P.SEIBLL MAURICE RABINOVIICH INVENTORS ATTORNEYCHJF SHEET PATENTED MAR l 8 I975 QAUQQQQQQQQ m HERVE 'E.BIBRING GEORGESP.SEIBEL MAURICE RABINOVITCH INVLNTORS BY WW! ATTORNLY PATENTED NARI81975 SHEET DSUF 1O PATENTEU MAR] 23 ms SHEET 06 0F 10 FIG. 9b

PATENTEU 3,871,835

SHEET 07 0F 10 FIG. IO

1 jig-1 lmg/c m rag/cm ig-l? 111mm @mumm I'IARICL lmBINOVIlCIi mmvronsATTORNEY PATENTEI] MARI 81975 SHEET lOUF 10 M l F REFRACTORY COMPOSITEALLOYS CONTAINING ROD-LIKE AND/OR PLATELET-LIKE LAMELLAE CROSS-REFERENCETO COPENDING APPLICATION The present application is acontinuation-in-part of our copending application Ser. No. 2,160 filedJan. 12, I970 (now abandoned).

FIELD OF THE INVENTION The present invention relates to refractorycomposite alloys containing single crystalline rod-like and/orplatelet-like lamellae, their manufacturing method, as well as anapparatus for carrying into effect said method.

BACKGROUND OF THE INVENTION The manufacturing of metallic compositecompounds, containing rod-like and/or platelet-like lamellae, byunidirectional solidification of eutectic binary alloys has already beenproposed, as for example in the KRAFT U.S. Pat. No. 3,124,452.

However, since the metallic matrix of a binary alloy is unable to offera sufficient resistance to hightemperature corrosion, only compoundshaving unsatisfactory properties and especially the inability to resisthigh-temperature corrosion have been prepared by such a method.

It has also been proposed to provide alloys with a ternary eutecticcomposition having two phases, one a matrix either of nickel, cobalt orchromium, the other a reinforcement phase which is oriented and in theform of fibers (fibrous phase) contributing high corrosion resistance atelevated temperatures. The fibrous phase is composed of a monocarbide ofa transition metal (i.e., titanium, zirconium, hafnium, vanadium,niobium or tantalum see US. Pat. No. 3,528,808).

There have also been suggested alloys consisting of a monovarientternary eutectic with a matrix of an alloy of nickel or cobalt and afibrous phase (phase ofa whisker or lamellar morphology in substantialalignement) constituted of a carbide or mixture of carbides oftransition metals (i.e., the carbides of niobium, tantalum, titanium,vanadium, zirconium, hafnium, chromium and cobalt); another prior-artalloy of the ternary type includes the system in which nickel aluminideforms the matrix and contains fibers of chromium (see U.S. Pat. No.3,564,940).

As has been pointed out in the literature, all of these alloys orsystems are of substantially monovarient composition, i.e., thedifference between the number of elements contained in the alloy and thenumber of phases is at most unity. This has posed severe restrictionsupon prior-art attempts to extend the principles described above to morecomplex systems and thereby to further improve the properties ofmetal-matrix/metal-carbide fiber systems. While we do not wish to bebound to any theory as to why the prior-art systems have been limited tomonovariance, as noted above, and have hitherto failed to comprehend avariance of two or three or more (i.e., systems in which the differencebetween the number of metallic elements present and the number of phasesis two, three or more), the reasons are probably the following:

the phase diagrams of quaternary, quinary or of higher-order systems arenot well known;

the necessity of taking into account the partition coefficients i.e.,the ratio, at a given temperature, of the concentration of an element inthe solid phase to its concentration in the liquid phase of all thepresent elements excessively complicates the establishment of thesolidification conditions to be used; and

in order to prepare materials are to be used at very high temperatures,melting and solidification must be achieved at very high temperatures,so that the operating conditions which must be used in theunidirectional solidification process become exceedingly difficult tomaintain.

The difficulties encountered with the aforementioned prior-art systemsare the following:

Alloys having a matrix based upon cobalt, nickel or iron manifest poorresistance to corrosion at high temperatures, especially when they donot contain a significant proportion of chromium. In such systems, thelimitation to monovariance or, as noted, the limitation upon the numberof chemical elements present because the final structure is obtained bydirectional solidification has the following consequences:

Where it is not possible to introduce chromium in sufficient proportionsinto the matrix, the resistance of the alloy to high-temperaturecorrosion is poor.

Where chromium is introduced into the matrix at higher levels, it isalso found-to be present in the reinforcement phase in the form of thecarbide Cr C As pointed out by J. R. LANE and N. 1. GRANT (Transactionsofthe A.S.M., XLIV, 1952, pp 113 137: Carbide Reactions in HighTemperature Alloys), at temperatures between 730 and 870 C, the lattercarbide tends to be transformed into Cr C with liberation of carbonwhich is retained in the matrix in the form, for example, of Cr C As aconsequence, a finely divided disperse phase is created. In effect, thefibrous phase of a reinforcement alloy provides strength-increasingfibers with a contribution to the tensile strength which depends uponthe continuity and homogeneity of the fibers, the latter being bondedtogether by the matrix. With carbon-loss transformation described above,both of these critical characteristics of the fibrous phase, as well asthe vital bond between the matrix and fibers, are adversely affected,the fibrous phase being replaced by a dispersed phase.

If one attempts to avoid this difficulty by increasing the chromiumcontent of the reinforcement phase, the partition function comes intoplay to simultaneously increase the chromium content of the matrixphase.

Since it is almost impossible to exclude trace elements, such asnitrogen, which render chromium alloys fragile when they include solidsolutions rich in chromium, this approach has been self-defeating.

OBJECTS OF THE INVENTION One of the main objects of the presentinvention is that of providing a new metallic composite product which iscomposed of at least four chemical elements, containing rod-like and/orplatelet-like lamellae imbedded in a metallic matrix, prepared byunidirectional solidification and which evidences good mechanicalproperties (especially flow stress) at room temperature as well as athigh temperatures.

One of the more particular objects of our present invention resides inproviding such a composite metallic compound which can stand drycorrosion at temperatures of about l,0O C, and which consequently is ofconsiderable technical interest.

Another object of our present invention is to provide ductile alloyswhich can be used in making pieces having significant deformationcapacity before breaking.

Another object of our invention is to provide alloys showing a highresistance to thermal strain, i.e., to repeated temperature variations.

Another object of out invention is to provide such alloys exhibiting ahigh yield point at high temperatures as well as at room temperature.

One of the principal objects of the present invention is, therefore, toprovide alloy systems with a reinforcement phase by directionalsolidification in which the degree of variance can exceed unity, therebypermitting optimalization of both the matrix and reinforcement-phaseelemental compositions.

Still another highly important object of the invention is to providealloys with high resistance to corrosion at high temperatures and yetfree from the increased fragility which has characterized the alloysproduced in earlier attempts to provide reinforcement-phase systems.

A crucial object of the invention is to provide an improved alloy withresistance to corrosion at elevated temperatures which possesses areinforcement phase which is absolutely stable at all temperatures up tothe temperatures of use.

It is also an important object of the invention to provide an orientedalloy comprising two or more reinforcement phases, and a process formaking the alloy, in which one phase is an aligned fibrous reinforcementphase in the matrix and the other reinforcement phase is dispersed inthe matrix and serves to increase the resistance of the alloy to shearand the alloy hardness.

Another object of the invention is to provide an ironbased, cobalt-basedor nickel-based matrix in an alloy which, in addition, compriseschromium and another metallic element such that the oriented (fibrous)reinforcement phase consists entirely of carbides of this other elementand is entirely free of chromium, i.e., there is no chromium carbide inor as the fibrous oriented reinforcement phase.

SUMMARY OF THE INVENTION The invention is based on our finding of theexistence, in the ternary CoTaC system, of a eutectic point in apseudo-binary diagram in which one of the phases is tantalum carbide,TaC, cobalt being the other phase. The pseudo-binary eutectic point,evidenced by thermal analysis, corresponds to a composition ofapproximately 13 percent by weight of TaC and to a temperature of aboutl,400 C; it is then possible to prepare compounds having excellentmechanical properties, by unidirectional solidification.

The invention is also based on the observation that the CoCr phasediagram has a very flat shape for chromium concentrations up to 25percent, and the application of our method to the complex system Co-CrTaC leads during the progressive solidification to a crystallizationof carbide TaC is a cobaltchromium solid-solution matrix. We have thusprepared a new composite metallic compound containing rod-like and/orplateletlike lamellae of TaC in a chromium-containing matrix, whichprovides oxidation resistance at high temperatures.

In its most general sense, therefore, the invention comprises an alloywith plural variance consisting of a matrix and a fibrous, orientedreinforcement phase wherein the matrix is essentially composed ofchromium and at least one other metal selected from Group VIII ofthe 4thPeriod of the Mendeleiv Chart ofthe elements, namely iron, cobalt ornickel, and the reinforcement phases consists essentially ofmonocrystalline particles of a carbide of a metallic componentsubstantially free from chromium and including at least one metalselected from Groups Nb and Vb of the Mendeleiev Chart.

The invention thus comprises the remelting- Iresolidification treatmentofa bar of an alloy of at least one metal of Group VIII of Period 4 ofthe Mendeleiev Chart of the elements, of chromium, of carbon and of atleast one of the transition metals of Groups Nb and Vb of the MendeleievChart, the transition metal aand carbon being present in the alloy instoichiometric proportions, under conditions which form a reinforcementphase (directed solidification). The treatment, where the chromiumcontent does not exceed 25 percent by weight, yields a matrix consistingof a solid solution of chromium in the metal of Group VIII, Period 4,and an oriented reinforcement phase constituted by monocrystallineelongated particles of one or more carbides of the metals of Groups Nband Vb. Thus the chromium is practically totally in solid solution inthe metal or metals of the matrix and the metal or metals of Groups Nband Vb are practically totally in combination with carbon in thereinforcement phase. There is, therefore, no chromium (Group Vlb) in thelatter phase to adversely affect the properties of the alloy.

Our method makes also possible the preparation by unidirectionalsolidification of new composite metallic compounds in which the cobaltpart is completely or partially played by nickel or iron, the inventionmaking use of the fact that the NiCr and Fe-Cr phase diagrams have alsoa very flat shape up to a relatively high ponderal composition ofchromium, of about 45 percent for the nickel-chromium diagram.

The new composite metallic compounds embodying our invention containthen rod-like or platelet-like TaC lamellae in a NiCr matrix, or in aCoNi-Cr matrix, or also in an FeNiCr matrix.

According to another feature of our invention, the matrix isadvantageously formed by an austenitic solid solution, e.g. one based oniron and/or nickel and/or cobalt, which crystallizes in theface-centered-cubic lattice.

Such a matrix does not show any allotropic transformation from roomtemperature to the melting point, and thus offers the advantage, amongothers, of not exhibiting any brutal dimensional variations, ontemperature changes, which could cause, for instance, more or lesslocalized deformation, or internal stresses.

According to one embodiment of the present invention, the matrix is madeof an NiCrAl alloy hardenable by precipitation.

The tantalum and hafnium carbides being miscible in any proportions, theapplication of our method to complex starting systems, including bothtantalum and hafnium, leads too new composite metallic compounds inwhich the rod-like and/or platelet-like lamellae imbedded in the matrixare made of a solid solution of tantalum carbide, TaC, and hafniumcarbide, l-lfC.

In many cases, tantalum can be replaced by other metals, especiallytransition metals of the Nb of the Mendeleiev Chart, i.e., titanium andzirconium, as well as transition metals of the Vb period, i.e.,vanadium, niobium and hafnium.

According to another feature of our invention, the metal or metals ofthe carbides of the unidirectional solidification or reinforcement phaseare incorporated in excess in the alloy so that one part of this metalor these metals forms with the constituent elements of the matrix.either a solid solution or one or more intermetallic compounds whichincrease the mechanical properties of said matrix, that part acting asan additional element for the matrix.

According to one preferred embodiment, titanium is the metal forming thecarbide of the reinforcement phase.

The expansion coefficient of titanium carbide is, indeed. very close tothat of the nickel-based and/or cobait-based matrices; this avoids, ontemperature variations, the development of any stresses along the matrixreinforcement-phase interfaces.

Nonrestrictive examples of complex starting systems, leading to alloysaccording to our invention are cited hereunder:

Co Cr-Ta-C Fe-Ni-Cr-Ta-C Co-Cr-Nb-C Ni-Cr-Ta-C Lo-Cr-Ta-Hf-C Ni-Cr-NbCCo Ni-Cr-Ta-C Ni-Cr-Al-Ta-C Co Ni-Cr-Nb-C Ni-CR-Al-Ti-C Co Ni-CnTa-Hf-CThe alloys according to our invention are made by progressiveunidirectional solidification in an apparatus which permitsestablishment of predetermined values of the operating parameters, suchas the temperature of the liquefied zone of a previously prepared rod,and of the resolidified zone, the movement rate of a solidificationinterface, the flat shape of a solidification interface. the temperaturegradients across said solidification interface etc.

Where the metallic component of the reinforcement phase (carbide-formingelements) are tantalum, titanium or niobium as noted above, theseelements are present in the alloy in amounts of substantially to 18percent by weight, 8 to 12 percent by weight and 7 to 17 percent byweight, respectively. The chromium content of the alloy is preferably 5to 25 percent by weight and the matrix phase is preferably an austeniticsolid solution. Furthermore, when aluminum is present as an additionalmetal other than the metals of Group VIII, Period 4 of the MendeleievChart and other than the metallic component of the carbide reinforcementphase. it is present preferably as an intermetallic compound with one ofthe other matrix-phase elements in an amount of substantially 2 to 6percent of the alloy. In another preferred embodiment of the inventionthe reinforcement phase is a single-phase solid solution ofsubstantially 80 percent by weight tantalum carbide and percent byweight hafnium carbide. The elon gated monocrystalline particles of thereinforcement phase preferably have a smallest transverse dimension ofsubstantially 0.3 to 2 microns and a length at least equal to 3.000times their smallest transverse dimension.

In other general terms, the invention also comprises an alloy of achromium-containing solid solution matrix and, imbedded in this matrix,a -directional (elongated) reinforcement phase in the form ofmonocrystalline particles (preferably extending parallel to each other)and a reinforcement phase in the form of substantially point-likeparticles dispersed in the matrix. As a divariant or polyvariant systemthe alloy can be considered to comprise at least four chemical elementsforming a metallic matrix (preferably containing chromium in solidsolution) and a directional reinforcement phase (chromium-free) in theform of elongated monocrystalline particles.

6. DESCRIPTION OF THE DRAWING Other features and advantages of ourinvention will be clear from the following description given withreference to the accompanying drawing in which:

FIG. 1 is a schematic view of an embodiment of an apparatus used inmaking alloys of the present invention;

FIG. 2 is a schematic view of a modified apparatus;

FIG. 3 is a diagram showing the temperature variation along a rod; I

FIG. 4 is a schematic illustration of another type of apparatus used inmaking alloys of the present invention;

FIGS. 5, 6 and 7 are photomicrographs of alloys of the presentinvention;

FIG. 8 is a diagram;

FIG. 9a, 9b, 10 13 are photomicrographs of alloys of the presentinvention;

FIG. 14 is a diagram;

FIGS. 15 and 16 are photomicrographs of alloys of the invention;

FIG. 17 is a diagram; and

FIG. 18 is a photomicrograph of an alloy according to the invention.

SPECIFIC DESCRIPTION Reference is first made to FIG. 1 which showsschematically an apparatus used in making the compositions according tothe present invention and which is of the electron-bombardmentfloating-zone type.

The rod 1 of the alloy to be treated in set up in vertical position, ina vacuum chamber (not shown) by clamping its two ends in locking means 2and 3, the latter being grounded. An annular electron gun 4 surroundsconcentrically the rod 1 and can be moved at constant rate in a verticaltranslation motion by means of a screw-and-nut drive mechanism (notshown).

The electron gun comprises a tantalum filament 5 heated to hightemperature by Joule effect and emitting electrons. The filament has ahigh negative voltage with respect to the grounded rod 1 which is to betreated. In order to focus the electrons which are bombarding the rodalong and annular zone, a tantalum chamber 6 placed around the filamentand rod has the same negative voltage as the filament. This arrangement,in which filament and melted zone 7 are not in a direct line of sight ofeach other, avoids any reciprocal contamination by the metallic vaporswhich both emit. The necessary thermal gradient is established bycooling the solidified zone by radiation.

At the beginning of the run,.the mounting of gun 4 is placed at thelower part of the rod; the rate of upward motion can range from 0.5 to30 cm/hour.

In FIG. 2 there is shown another apparatus used in producing the alloysaccording to our invention. The

apparatus comprises a fixed refractory tube 10 inside which a metallicrod 11 is placed, the composition of the rod corresponding to that ofthe desired alloy. The tube 10 is surrounded by another fixed tube 12 oflarger diameter, the space 13 between the two tubes being filled with aprotective gas circulating from an inlet pipe 15 to an outlet pipe 15.The tubes 10 and 12 are coaxial and their common axis is vertical.

The tube 12 is surrounded by a resistance furnace 16 which here consistsof three superposed stages 17, 18 and 19. Each stage comprises a heatingresistance 20, 21 and 22, respectively generating temperatures of atleast l.500 C. Each stage has its own regulation means. The resistancesare surrounded by respective refractory masses or rings 23, 24 and 25.The furnace 16 has a length equal to at least 10 times the diameter oftube 10.

The furnace is continued in its lower part by a cooling device 26, whichcomprises a body 27 with a flange 28, and surrounding that body a coil29 in which circulates a cooling fluid.

The unit including the furnace l6 and the cooling device 26 is assembledin such a way that it can be moved, by means of a suitable drivemechanism (not shown), with respect to tubes 10 and 12 along atranslation path parallel to the common axis of the tubes. The tube 10is placed in a holder 30, cooled by circulating fluid as schematicallyrepresented by the two sections 31 and 32 on opposite sides of apartition element 33.

The operation is as follows:

The stages l7, l8, 19 of the furnace are adjusted in order to provide atemperature distribution along the axis of rod 11 as representedschematically by the diagram shown in FIG. 3.

On the left side of that FIGURE the rod 11 is represented in the sameposition as in FIG. 2. The arrow indicates the movement of rod 11 withrespect to the heating stages 17, 18 and 19. On the right side of thediagram the curve D represents the temperature variation along the rodfor a given position of the latter. In this diagram, the r axis is thetemperature axis and the 1 axis represents the abscissae ofthe rodsections. Temperature I] is the melting and solidification point of thematerial of the rod.

The temperature curve D can be divided into three sections:

a section 0 having a steep slope, positive with respect to the 1 axis;

a section b having a weak positive slope a section c having a very steepnegative slope.

The intersection points of the curve D with the line passing through thepoint t, and paralleling the abscissa determine on the rod the limits ofthe unmelted parts 5, of the rod. of the liquid part L and of theresolidified part S.

The particular shape. hereabove defined, of the temperature curve isobtained by giving to the melted zone L a large length (at least equalto the diameter of the rod) and operating the lower heating stage 17 athigher heating power than the other stages, taking into considerationthe influence of the cooling devices.

A very steep tmeperature gradient is thus obtained across thesolidification interface f, at the common boundary of the parts L and Sof the rod. This gradient ranges from about to 150 C per cm.

The apparatus hereinabove described also provides a prefect flat shapefor said interface, so that the microstructure of the solidified part Sis accurately directed oriented parallel to the axis of the rod.

The unavoidable variations of the translation rate or of the heatingpower have no detrimental influence on the structure formed at thesolidification interface. essentially because of the considerable massof liquid matter placed above said interface. Thus, band" formationwhich often occurs in directed single-crystal growth by moving asolidification interface is avoided.

Reference is now made to HO. 4 showing another type of apparatus used inmaking metallic alloys according to the invention. Induction heating isused: the induction coil 40 surrounds an other tube 41 made of drawnquartz for example. Heating is achieved by means of a graphite tube 42,acting as a susceptor, which heats by radiation a refractory tube 43 inwhich the rod 11 to be treated is placed. The tube 43 is placed on aholder 44 which also acts as a cooler; its body is composed of awater-cooled solid-copper cylinder. Water is admitted through a centralpipe 47 and then flows into an annular space 48 between pipe 47 and body44.

The lower end ofthe outer tube 41 is cooled by water circulation in achannel 45.

An inert gas is admitted through an inlet 49 to the annular space 50,existing between the surrounding tube 41 and the graphite tube 42, andis exhausted through an upper outlet 51.

The unit including the furnace and the cooling device is verticallyplaced and enables movement of the rod with respect to the inductioncoil, by means of a suitable drive mechanism (not shown).

Good results have been obtained with a 250 mm long susceptor having a 16to l8 mm diameter, surrounding a refractory tube of 8 to 12 mm diametermade of the alloy to be melted. The protective inert gas had a flow rateof 0.5 liter/minute. The constant translation rate of the beamsupporting the assembly ranged from 1.18 to 30 cm/hour. Thehigh-frequency generator had a frequency of MHz. The induction coil hadnine turns which were 1 cm apart. The regulation device enables atemperature control with a better than 1C accuracy.

8. SPECIFIC EXAMPLES The following are examples of alloys according tothe invention.

EXAMPLE 1 The workpiece initially consists of a quaternary castalloywhose composition (sgt percent) is as follows:

The workpiece is zoned by means of the apparatus shown in H6. 4. Therelative speed V of the rod is fixed at 1.2 cm/hour. The induction inputis adjusted so that the length of the liquid zone is equal to or largerthan five times the diameter of the rod, which is here 8 mm. For largerdiameters, this ratio can be reduced to 1. Under these conditions, thethermal gradient normal to the solidification interface is about C/cm.

Physico-chemical studies show that the final product is a matrix ofcobalt-chromium solid solution, inside which are imbeddedmonocrystalline thread-like particles of tantalum carbide, TaC,perfectly lined up parallel to the rod axis. These particles arelamellae having a more or less regular polygonal shape.

The ratio of the lamellae to the cylinders depends on the operatingconditions. In any event the smallest dimension of these particlesranges between 0.3 and 2 microns and their length is more than 3,000times their smallest dimension.

The nature and the structure of the particles, on the one hand, and thecomposition of the solid-solution matrix, on the other hand, give thematerial an excellent tensile strength at high temperatures (about 40hbars at l,000 C).

The breaking time, under flow, at l,000 C and under a stress of 10.5hbar is 700 hours in air. At l,050 C under ll hbar, the breaking timeunder vacuum is 2,500 hours.

According to the properties intended for the material, the chromiumconcentration of the solution and also, to a certain extent, the carbideconcentration of the rod can be varied to a large extent.

Very good results are obtained because of the plateau" shape of theCo-Cr phase diagram for high chromium concentration as will be seen inthe following examples.

EXAMPLE 2 The workpiece has the following starting composition (wgtpercent):

The operation is carried out as indicated in Example 1 hereinabove, butwith a movement rate V of 2.5 cm/h. The final material has excellentmechanical properties at high temperatures as well as at roomtemperature. On tensile tests, at room temperature, the yield point is100 hbar, the ultimate tensile strength R is l hbar and the elongationat rupture is percent (FIG. 8). The fatigue strength at 10 cycles inrotative bending is i 63 hbar.

At high temperatures, such a material exhibits a good mechanicalstrength and corrosion resistance.

On flow tests in air, the breaking time at l,000 C under 10.5 hbar is1,850 hours.

70 hbar 40 hbar.

At 800C 2 R and at l000C I R Such a material has also at roomtemperature, besides a high mechanical strength, a very good ductility.

This latter property is probably due to the fact that the structureofthe reinforcement phase lets it play entirely its part, said phaseoccupying a rather small volume of the matrix which consequently impartsto a piece made of such an alloy a substantial deformation capacitybefore breaking.

The operating conditions of the furnace can be moditied to quite a largeextent. Thus the movement rate of the solidification interface can beadjusted to a value rangmg from 1 to 15.5 cm/hour.

As for the thermal gradient, normal to the solidification interface, itsvalue can be lowered to 30 C/cm in certain cases, with, stillinteresting results.

The chromium concentration of the matrix, the movement rate of thesolidification interface and the thermal gradient at the interface canbe adjusted in order to have the reinforcement phase consist almostexclusively of cylinders, or two-dimensional lamellae, ortriple-branched lamellae, as shown on photomicrographs of FIGS. 9a and9b (X 2900) which show the influence of an addition of a 15 percentamount of chromium upon the carbide whiskers. FIG. 9a is a scanningelectron photomicrograph taken after electrolytic selective etching ofthe matrix revealing the whiskers, taken at an angle, for an alloy whosecomposition is:

Cobalt 87 7: by weight Tantalum: l2.2 7r by weight Carbon 0.8 '7: byweight (The latter alloy, for comparison, was unidirectionally solidifidby zoning in the apparatus of FIG. 1 at a movement rate V of thesolidification interface of 5.25 cm/hour and a thermal gradient G at theinterface of about 500 C/cm, the length of the liquid zone beingapproximately equal to the bar diameter. A photomicrograph of thecross-section is shown in FIG. 5 (X 780); when the rate was increased toV 15.75 cm/hour the cross-section after electrolytic etching had thephotomicrograph (X 4,000 with scanning electron microscope) of FIGS. 6).

FIG. 9b is a similar photomicrograph for an alloy whose compositioncorresponds to that of Example 2 with a rate V adjusted to 1.2 cm/hour.

EXAMPLE 3 The operation is carried out as in the preceding Example, thecomposition being as follows (weight percent) By unidirectionalsolidification at a rate V adjusted to 1.15 cm/hour, a final product isprepared which is EXAMPLE 4 The workpiece consists originally of aquaternary alloy having the following composition (weight percent):

This workpiece is zoned by the unidirectionalsolidification method, theinterface moving at the rate of 1.2 cm/hour.

An alloy is obtained whose structure is shown on the photomicrograph ofFIG. 11 (X 3,000), taken at an angle after selective electrolyticaletching, and whose reinforcement phase whiskers have approximatelysquare sections. The alloy is a nickel-chromiumtantalum solid-solutionmatrix in which are imbedded long-sized monocrystalline particles oftantalum carbide. The matrix, which includes the excess of tantalum notcombined with carbon into tantalum carbide, has a higher hardness than anickel-chromium solid solution having proportions of nickel and chromiumsimilar to those of matrix consisting of only these two components.

In the following Example the nickel-solid-solutionbased matrix isreinforced by niobium carbide:

EXAMPLE The workpiece has the following starting composition (weightpercent):

After zoning as described in the foregoing Example, a compound isobtained whose long-sized carbide whiskers also have a cylindricalshape, as shown in FIG. 12 (X 5.500) which is an electronphotomicrograph taken at an angle of40 with respect to the axis of therod and with a scanning microscope after partial extrication of thewhiskers by electrolytic selective etching ofthe matrix.

Thc thus-prepared alloy is a matrix whose hardness is higher than thatof a nickelchromium solid solution having proportions of nickel andchromium similar to those of a solid solution consisting of only thesetwo components.

EXAMPLE 6 The workpiece has the following starting composition (weightpercent):

By the unidirectional-solidification method described above. an alloy isprepared in the form of a nickelchromium matrix in which long-sizedmonocrystalline particles of titanium carbide are imbedded.

The properties of this alloy are particularly remarkable because on theone hand the titanium carbide. whose expansion coefficient is close tothat of the matrix, forms a reinforcement phase which prevents stressformation at the matrix reinforcement-phase interfaces, and. on theother hand the excess titanium forms a dispersed compound Ni Ti whichincreases the matrix hardness and whose dispersion can be improved by astructural precipitation treatment.

Although the introduction in excess of carbideforming metals in Examples4, 5 and 6 above leads to overall compositions which slightly deviatefrom eutectic compositions, the structures obtained remain of theeutectic type.

Because of the closely similar cobalt, nickel and iron properties, it isalso possible to prepare alloys wherein a transition-metal carbide formsa reinforcement phase imbedded in a solid-solution matrix, for instancea Ni- Co-Cr or Ni-Fe-Cr matrix.

The composition is advantageously so chosen, for such alloys, as toyield a matrix constituted by an austenitic solid solution.

Note that austenitic solid solution" taken in its most general sense,means a solid solution based on iron and- /or nickel and/or cobalt whichcrystallizes in the facecentered-cubic lattice.

EXAMPLE 7 The workpiece has the following starting composition (weightpercent):

The workpiece is unidirectionally solidified by means of the apparatusdescribed with reference to FIG. 4, at a solidification rate of 1.2cm/hour.

The final product is a CoNi-Cr solid-solution matrix, having aface-centered cubic structure, in which long-sized monocrystallinerod-like and/or platelet-like lamellae of tantalum carbide are imbedded.The matrix keeps its austenitic structure for temperature values rangingfrom room temperature to its melting point. The structure of the alloyis shown in FIG. 13 which is a scanning electron photomicrograph (X1,900) taken at an angle of 40.

The mechanical characteristics of such an alloy are the following:

11 I'OOI'TI temperature:

Besides, this alloy has a very good resistance to dry corrosion. asshown by thermogravimetry, indicating a weight increase of l mg/cm2after 17 hours heating at I.OOO C in air,

The weight increase follows a law as shown by the graph of FIG. 14 inwhich the ordinates indicate the weight increase in mg/cm2, time beingplotted on the abscissae.

Comparative tests have been conducted with the alloy of the inventionand other refractory known alloys such as those known as IN (cast) andUDI- MET 700 (forged).

i. Shock bending The tests have been conducted at 20 C, 700 C, and l,000C on cylindrical test pieces with a Wolpert Feston PWS hammer-machine.

The results of the tests are summarized in the following table A whichshows the advantages of the alloy according to the invention.

TABLE D Test temperature Time in hours (i) K=T(20+log l) C K Stress forelongation at for at (T) (hbar) rupture A% 0.5 rupture TABLE A 15 Thetwo last columns refer to the values of the Larson-Miller index Kdefined by K T (20 log I) 10' Alloy Temperature C Resilience da .l/cm2wherein T is the temperature in K and t the duration of the testmeasured in hours. The indicated values of Alloy of the 20 21 mentionthis index emphasise the good characteristics of the UDIMET 700 20 ll 20alloy according to the invention.

iN 100 20 6 Alloy of the 700 21 EXAMPLE 8 invention HL EE $88 2 Theworkpiece has the following starting composi- Auoy of he 1000 tion(weight percent):

invention 25 UDIMET 700 1000 20 IN 100 1000 2.5 CO Cr 1 20% ii. Thermalfatigue 10% Nb: 7.1% Wedge shaped test pieces are placed in the flame ofC 0.9%

an air-propane burner in order that the edge of the test piece bebrought to 1,100 C and stays at this temperature during 60 seconds. Thetest piece is then withdrawn from the flame and cooled in 20 seconds to300 C by a cold air stream.

A test piece made of the alloy according to the invention was subjectedto a 300 cycles test. The first fatigue crack appears after 100 cyclesand its propagation is relatively slow. The same test carried on goodrefractory nickcl base alloys such as [N 100 or PD 16 results in thefirst fatigue crack appearing after 10 to 20 cycles and the propagationis much quicker than with the alloy according to the invention.

Further tests have also been carried out namely fatigue test and flowstress test.

The results of these tests are summarized in the following tables: iii.fatigue test at room temperature the fatigue strength in rotaat 800 Cthe fatigue strength under ondulated traction (87 Hz) is as below:

TABLE C Stress mini maxi Number of cycles 2 40 10 2 45 rupture iii. flowstress test The results are summarized in Table D With the samedirectional-solidiiication method as described in the foregoing Example,an alloy is prepared whose structure, as shown in FIG. 15 which is anoptical microphotograph of a cross-section of the sample (X 780), is amatrix consisting of a face-centeredcubic Co-Ni-Cr solid solution inwhich longsized monocrystalline niobium carbide particles are imbedded.

At 1,000 C, the U.T.S. of such an alloy is 44.5 hbar.

The cobalt of the solid solution can also be replaced by iron, a metalbelonging like the former to group V111 of the fourth period of theMendeleiev Chart. A material displaying very interesting characteristicsas regards dry corrosion can be prepared as follows:

EXAMPLE 9 The workpiece has the following starting composition (weightpercent):

: 48.5% 22 Ni 17.5% Ta 11.25% C 0.75%

Fe Cr EXAMPLE 10 Operating as in the previous Examples, but replacingtantalum by hafnium (either totally or partiallylinasmuch as thecarbides of these two metals are miscible in any proportions, yields areinforcement phase of monocrystalline particles of mixed single-phasecarbides, TaCHfC.

The workpiece has the following starting composition (weight percent):

Co 57% Cr 20% Ni: l7 Ta: 9.7% Hf: 2.57: C 0.8%

The structure of the material prepared by the method just described, ata rate V of L2 cm/hour, is shown in the electron photomicrographrepresented in FIG. 18 (X 5.800).

It is to be noted that the crystalline-phased solid solutioncorresponding to 80 percent of TaC and 20 percent of HfC is at thepresent time the most refractory mixed crystal known (melting pointofabout 3.940C).

According to our invention. also. an additional element is introducedinto the alloys prepared as hereabove described. this element improvingthe mechanical properties of the matrix by a subsequent thermaltreatment or by precipitation on solidification. Thus in thenickel-based alloys aluminum is added in a proportion ranging from 2 to6 percent.

Examples of such alloys are as follows:

EXAMPLE 11 The workpiece has the following starting composition (weightpercent):

Ni Ta Cr C Al: 3

By unidirectional solidification, an alloy is prepared in which theexcess tantalum is dissolved in the matrix inside which the tantalumcarbide reinforcement phase is imbedded. and which comprises anintermetallic complex compound of the type. appearing on cooling aftersolidification; its dispersion can be improved by a structuralprecipitation treatment.

The structure of such a material is shown in FIG. 5 which is an electronmicrophotograph taken at an angle of 24 (x 5,200). It can be seen thatthe carbide whiskers having a square section are imbedded in the complexmatrix containing the hardening precipitate y. The photomicrograph showsthe structure after roughsolidification, but the latter can be refinedby thermal treatments.

Direct observation by transmission electron microscopy shows, asevidenced in FIG. 6 (X 27,000), the particular shape of the y-typeprecipitate and the dislocations existing in the matrix.

Before any treatment, this material as a U.T.S., at room temperature ofI25 hba (l78 X 10 psi) with 10 percent elongation.

EXAMPLE Ila The workpiece has the following starting composition (weightpercent):

Ni 53.6 7: Co 20 71 Ta ISI 7: Cr: l0 7? C 0.4 7: Al 3 Z 5 The alloyobtained in this example has a structure similar to that of Example 11but differs therefrom by the replacement of part of the nickel bycobalt. lmmel0 diately beneath the solidus this alloy is constituted byan austenitic matrix of Ni, Co, Cr, Al. into which are imbeddedreinforcement fibers of tantal carbide.

During the cooling a y phase of Ni--Al compound precipitates in thedispersed state. Dispersion can be improved by an homogenization thermaltreatment of 1 hour at 1,l50 C followed by a structural precipitationtreatment of 24 hours at 760 C.

The tests carried out on this alloy are summarized in the followingtables:

TABLE E 25 (at room temperature) solidification Thermal UTS Elongationat rate (cm/hour) treatment (hbar) rupture (71) 2.36 144 9.5 3.50 149.52.36 l l 146 12.8 3.50 (2) l46.5 9.2

2.36 (2 I 161 l l .2 3.50 (2) 166 4.1

Thermal treatment: (I); homogeneisationzl hour at ll50C (2 l lstructural precipitation treatment 24 hours at 760C TABLE F (atdifferent temperatures) Solidification Thermal Test UTS Elongation rate(cm/hour) Treatment temperature (hbar) at rupture 2.36 without 760 85 92.36 (2) 760 98.5 10.8 236 without 1000 35.2

2.36 without 1000 33.2 8.3 2.36 (2) 1000 4L3 4.45

Thermal treatment: l homogenizationzl hour at llC (2): l) structuralprecipitation treatment 24 hours at 760C. 50

TABLE G Thermal Test Stress Number Treatment Test Tempera- (hbar) ofture cycles without Rotative binding 20C 1 45 10 (49 HZ) do. do. do. i50 rupture (2) do. do. :45 l0 (2) do. do. i 50 rupture without Ondulatedtraction 800C mini 2, l0

(87 Hz) maxi 60 do. do. 800C mini 2. rupture maxi 65 Thermal treatment:l homogenizatioml hour at l l50C.

(2): l structural precipitation treatment 24 hours at 760C.

(Flow stress tests) fhermal fest Stress Time in hours (t) K=T(20+log 1)l Treatment Temperature (hbar) "C K) for to for to T elongation ruptureA 0.5% rupture without 760 1033 60 2 40 21 22.3 {2) 760 i033 60 96 76022.6 23.6 without 850 I123 40 6 8 23.4 23.5 (2) 850 1123 40 35 I42 24.224.8 without i000 1273 i 351 483 28.7 28.8 l2) i000 I273 43 I020 27.529.3 without I030 i303 i2 69 70 28.5 28.5 (2) i030 1303 i2 69 316 28.529.4 without 1070 1343 i2 14 15 28.5 28.4 l2) 1070 1343 i2 34 34 29 29Thermal treatment: (I): homogenizatioml hour at ll50C (2): l structuralprecipitation treatment 24 hours at 760C.

As defined In lable Dv EXAMPLE l2 (4, 5, 6, ll, 12 and 13), dealing witha nickel- The workpiece has the following starting composichrPmlum mamxcontam about 85% by weight of the matrix-forming elements and about 15non (welght percent).

percent by weight of carbide-forming elements so that 25 a maximum of 45percent by weight of chromium in the Ni 72 matrix phase corresponds toalloys containing up to 37 5 1 '8 percent by weight chromium. Al 3 2 Wehave used herein the terms rod-like and c 1 0.7 7 platelet-like lamellaeto describe the elongated monocrystalline particles constituting thereinforce- After Zoning, an alloy is obtained whose structure is mentphase and we intend the term monocrystalline, similar to that of theforegoing example, but in which therefore, to comprehend both fibrousand flattened tantalum is replaced by niobium. reinforcement particles.

The alloys of the present invention have proven to be EXAMPLE 13especially effective for use in high-temperature gas tur- The workpiecehas the following starting composig i ali'ronauncalh.agphcatlons m termsof on g p c anic a res stance to lg temperatures, resistance to materialfat1gue, high corrosion resistance and exceptional mechanical strengthon a long-term basis. Ni: 76.5 7. 40 We claim Ti l. A refractorydirectionally solidified polyvariant fi- $1 9 bet-reinforced compositehaving eutectic-type struc- Al 4 ture consisting essentially of twodistinct independent phases constituted by: After treatment by thedescribed method, at a rela- 5 a. a complex multicomponent matrix phaseconsisttive speed of 2.5.cm/hour, an alloy is obtained which ingessentially of: contains a titanium carbide reinforcement phase, imi. atleast one metal selected from the group conbedded in a nickel-chromiummatrix comprising an insisting of Fe, Ni and Co, and ter-metalliccompound Ni (Ti,Al), which has a faceii. chromium in an amount between10 and 25 percentered-cubic structure, and whose dispersion can be centby weight of the composite; improved by a structural precipitationtreatment. FIG. and in said matrix: 7, which is a scanning electronphotomicrograph (X b. an in situ grown reinforcing phase free from chro-2.300) taken at an angle of 40, shows the structure of mium andconsisting essentially of whisker-like ihe reinforcement PhaseContaining finely dispersed elongated monocrystalline fibers of at leastone point-like carbide particles, and also the y-type p ipmetalmonocarbide, the metal of which is selected itate as it appears on roughsolidification. from the group constituted by Ta, Nb, Hf and T1. While acritical characteristic of the alloys of the 2, A refractorydirectionally solidified composite as present invention is the exclusionof chromium in the d fi d i l i 1 h i id metal f h monocarreinforcementphase, it should be observed that up to bid i l d s tantalum in anamount of 10 to 18 per- 25 percent by weight of chromium may be presentin 0 t by weight of the composite. solid solution in the matrix phasewhere the latter is pri- 3. A refractory directionally solidifiedcomposite as m rily a cobalt-based system and up to 45 percent bydefined in claim 1 wherein said metal of the monocarweight in solidsolution in the matrix phase where the bide includes titanium in anamount of 8 to 12 percent latter is primarily a nickel-based system.Where mixby weight of the composite. tures of cobalt and nickelconstitute the matrix phase, 5 4. A refractory directionally solidifiedcomposite as the chromium may be present in an amount of 25 percent byweight of the cobalt 45 percent by weight of the nickel. as a maximum.in the foregoing Examples defined in claim 1 wherein said metal of themonocarbide includes niobium in an amount of 7 to 17 percent byweight ofthe composite.

1. A REFRACTORY DIRECTIONALLY SOLIDIFIED POLYVARIANT FIBERREINFORCEDCOMPOSITE HAVING EUTECTIC-TYPE STRUCTURE CONSISTING ESSENTIALLY OF TWODISTINCT INDEPENDENT PHASES CONSTITUTED BY: A. A COMPLEX MULTICOMPONENTMATRIX PHASE CONSISTING ESSENTIALLY OF: I. AT LEAST ONE METAL SELECTEDFROM THE GROUP CONSISTING OF FE, NI AND CO, AND II. CHROMIUM IN ANAMOUNT BETWEEN 10 AND 25 PERCENT BY WEIGHT OF THE COMPOSITE; AND IN SAIDMATRIX: B. AN IN SITU GROWN REINFORCING PHASE FREE FROM CHROMIUM ANCCOSSISTING ESSENTIALLY OF WHISKER-LIKE ELONGATED MONOCRYSTALLINE FIBERSOF AT LEAST ONE METAL MONOCARBIDE, THE METAL OF WHICH IS SELECTED FROMTHE GROUP CONSTITUTED BY TA, NB, HF AND TI.
 2. A refractorydirectionally solidified composite as defined in claim 1 wherein saidmetal of the monocarbide includes tantalum in an amount of 10 to 18percent by weight of the composite.
 3. A refractory directionallysolidified composite as defined in claim 1 wherein said metal of themonocarbide includes titanium in an amount of 8 to 12 percent by weightof the composite.
 4. A refractory directionally solidified composite asdefined in claim 1 wherein said metal of the monocarbide includesniobium in an amount of 7 to 17 percent by weight of the composite.
 5. Arefractory directionally solidified composite as defined in claim 1wherein said matrix phase is an austenitic solid solution.
 6. Arefractory directionally solidified composite as defined in claim 5wherein said matrix phase is made of cobalt, chromium and up to 10percent by weight of the composite of nickel or iron, whereby anydetrimental allotropic transformation of the Co base matrix is avoided.7. A refractory directionally solidified composite as defined in claim 5wherein said matrix phase is made of iron, chromium and up to about 20percent by weight of the composite of nickel.
 8. A refractorydirectionally solidified composite as defined in claim 5 wherein saidmatrix phase is made of nickel, chromium and up to 20 percent by weightof the composite of cobalt, whereby a greater volume fraction of fibersis allowed.
 9. A refractory directionally solidified composite asdefined in claim 1 wherein said matrix phase includes at least onemetallic addition element other than Fe, Ni, Co and those forming saidmonocarbide, said addition element forming a second reinforcing phaseconsisting of an intermetallic compound with a metal of the matrix phaseby soLid state thermal treatment.
 10. A refractory directionallysolidified composite as defined in claim 9 wherein said addition elementis aluminum in an amount of substantially 2 to 6 percent by weight ofthe composite.
 11. A refractory directionally solidified composite asdefined in claim 9, wherein said second reinforcing phase is in the formof substantially point-like particles dispersed in said matrix.
 12. Arefractory directionally solidified composite as defined in claim 8wherein said matrix phase further comprises a gamma '' type precipitateof nickel and aluminum.
 13. A refractory directionally solidifiedcomposite as defined in claim 6 wherein said matrix phase contains alsoat least one of the monocarbide forming metals.
 14. A refractorydirectionally solidified composite as defined in claim 8 wherein saidmatrix phase contains also at least one of the monocarbide formingmetals.
 15. A refractory directionally solidified composite as definedin claim 1 wherein said reinforcing phase is a single-phased solidsolution consisting of 80 percent by weight tantalum carbide and 20percent by weight hafnium carbide.
 16. A refractory directionallysolidified composite as defined in claim 1 wherein said fibers have asmallest transverse dimension of substantially 0.3 to 2 microns and alength at least equal to 3,000 times their smallest transversedimension.